Apparatus and method for subsurface structural modification of materials at reduced temperatures

ABSTRACT

Nanostructured or ultra-fine grained metallic systems according to embodiments of the invention may be formed of: pure Cu, pure Fe, or pure Ti, with grain sizes of less than 140 nm, 348 nm, or 59 nm, respectively. The metallic systems demonstrate a monotonically increasing grain size dependence from a minimum value attained at the surface; and a converse relation of microhardness, decreasing from 160 kg/mm 2 , 265 kg/mm 2 , or 320 kg/mm 2 , respectively. The grain refinement process at cryogenic conditions relies on the suppression of room temperature dislocation-mediated deformation mechanisms which facilitate grain restructuring, relaxation, and reorientation. At the cryogenic conditions, alternative mechanism for grain refinement, such as shear localization or dynamic recrystallization may be more dominant. Processes for refining the grain size of these metallic systems may include: subjecting metal plates to a high-energy milling process using a high-energy milling device to impart high impact energies to its surface. Due to the high-efficiency of this attrition process, these metallic systems are ideal candidates for improved corrosion and wear resistance.

CROSS-REFERENCE TO RELATED APPLICATIONS

This application claims priority to U.S. Provisional Application No.62/068,994, titled “Apparatus and Method for Subsurface StructuralModification of Materials at Cryogenic Temperatures,” filed on Oct. 27,2014 which is hereby incorporated by reference herein including allattachments and papers filed with U.S. Provisional Application No.62/068,994.

GOVERNMENT INTEREST

The invention described herein may be manufactured, used, and licensedby or for the United States Government without the payment of royaltiesthereon.

BACKGROUND

1. Field of the Invention

The present disclosure relates to nanostructured surface coatings andsubsurface modifications, having uniform to gradient microstructures,with single or multiple elemental constituents, and methods of makingthe same.

2. Background of the Invention

Nanostructured and nanocrystalline materials are well known to havenovel and/or enhanced properties compared to their coarse-grainedcounterparts. Specifically, nanostructured metals may have enhancedstrength, hardness, and/or improved workability. It has been shown thatthese nanostructures, when formed as metallic surface layers, couldresult in a number of significant improvements and new applications. Forexample, failure in metallic materials, such as fatigue fracture,fretting fatigue, or wear and corrosion often initiate and propagatefrom the surface of the material. As such, optimization bynanostructuring of the surface layer through surface modification wouldlimit such detrimental behavior and, in turn, not only enhance thesurface, but also the global or bulk properties of the material.

However, the abrupt transition in physical characteristics from such asurface layer to the bulk would likely result in a correspondinglydifferent response to physical phenomena as well. Therefore, it ishypothesized that creation of a gradual transition region between theunaltered bulk and this surface coating, having a range in properties,would be more effective in mitigating such phenomena.

There are a number of techniques to generate nanostructured surfaces byusing traditional coating techniques, e.g., plasma vapor deposition orchemical vapor deposition.

In the recent decade, similar to the well established method of shotpeening, a novel Surface Mechanical Attrition Treatment (SMAT)technology has been successfully applied to many metal-based materialsystems such as those based on Ni, Fe, Mg, Al, or Ti.

Surface mechanical attrition treatments have received much attention dueto their ability to enhance physical properties, such as yield strength,hardness, wear, and fatigue resistance in structural parts.Fundamentally, such treatments are carried out by repeated impacts ofthe metal surface with some impingement media which normally occurs athigh strain rates inside a closed or partially open container. Partiallyinelastic collisions during processing transfer kinetic energy into thesubstrate in the form of deformation strain energy and heat. In theinitial stages of SMATing it is expected that, at the surface, localizeddeformation occurs through compressive stresses and shear bandingfollowed by heavy deformation, wherein the total strain energy increasesdue to the creation of excess dislocations, stacking faults, and othersubgrain structures leading to an eventual attenuation and reduction ofthe grain size. Finally, an equilibrium grain size is reached, thespecific length scale of which is determined by a balance between theflux of induced defects, the evolution of the generated defect cellstructures, and their relaxation, reorganization, and restructuring. Asrelaxation/restructuring is thermally activated, the local as well asglobal temperature rise during processing plays a significant role indetermining the resultant microstructure during SMAT processing.

The prior art teaches the conventional application of SMAT, wherein,plastic deformation, at high strains and strain rates, is imposed on thesurface layers of bulk samples. Typically, spherical steel (or ceramic)balls of a few millimeters in diameter with smooth surfaces are placedin the bottom of a chamber that is sonically vibrated (usually with afrequency of 50 Hz˜20 kHz). When the balls are resonated, the samplesurface is impacted by a large number of flying balls over a shortperiod of time. Each impact of a flying ball (with a velocity of 1˜20m/s) induces plastic deformation at a high strain rate in the surfacelayer. Repeated multidirectional impacts result in repeated plasticdeformation in the top surface layer and the creation of bulk structuraldefects (dislocations, subgrain and stacking fault formation, etc.)followed by a process of grain restructuring, relaxation, andreorientation, which, in turn, induces progressive grain refinement downto the nanometer regime. What differentiates this form of SMAT from shotpeening is that in the latter case, the balls are directed onto thesurface using a nozzle. As such, the kinetic energy and directionalityof each ball is roughly the same as another ejected from the nozzle,rendering no major variation in the total quanta of deformation energybeing imparted to the surface.

Surface mechanical attrition treatment offers some unique advantagesover coating and deposition methods. For one, there is no change inchemical composition from the surface to the bulk or matrix. Anotheradvantage is the relative ease in the ability to create a continuous,gradient variation from the top surface, which is nanocrystalline, lessthan 100 nanometers, to the sub-surface, which is ultra-fine grained,less than one micrometer, to the bulk, which is coarse-grained, greaterthan one micrometer. With SMAT, bonding between the surface and matrixis not an issue.

The observed surface attrition process presents similar to that ofmechanical alloying/milling during active communication where theimparted kinetic energy depends on the frequency and the amplitude ofthe specific mill used; and, like in mechanical milling, a similaranalogous temperature rise (50-150° C.) in the substrate is expected tooccur for higher energy processes. This is primarily due to thedeformation induced heat generated within the substrate, ball to ball,ball to wall collisions, and frictional heating effects during therepeated attrition process. While generally this temperature rise isconsidered small, it can have a noticeable effect on certain metals andalloys (e.g., high purity, face centered cubic [FCC] nanocrystallinemetals which have been shown to undergo grain growth at roomtemperature).

Here we describe the procedures and advantageous effects of cryogenicSMAT processing at sub-ambient temperatures, preferably below theductile-brittle transition temperature of the material being treated,more preferably, below −50° C., most preferably, at liquid nitrogentemperature for commercially available oxygen free high conductivitycopper (Cu) (OFHC), pure iron (Fe), and pure titanium (Ti), eachrepresentative of a different crystallographic system. In the spirit ofthis invention, these exemplary systems are used as illustrations forwhat is possible with our concept. However, it is noted and emphasizedthat the applicability of the processes described herein have greaterutility to a much broader range of material systems, including, but notlimited to, other pure metals and alloys with commercializationpotential. In turn, to demonstrate the unique, non-obvious, and salientfeatures of this invention, we will differentiate the low temperaturestructural evolution of these metals from that occurring during roomtemperature processing.

Note, while SMAT processes have been applied to a number of systemsincluding, essentially pure metals such as Cu, Ni, Ti, Fe, and alloyssuch as stainless steels, currently, there is no or very limited data ondecoupling the thermomechanical effect during the processing of suchmaterials, especially at low temperatures where dynamic recovery andrecrystallization are dramatically suppressed.

This invention utilizes equipment, developed for high energy cryogenicmechanical alloying/milling, which presents a unique, non-obviousopportunity when applied to low temperature or cryogenic SMATing.

BRIEF SUMMARY OF THE INVENTION

Various unary, binary, and higher order metallic systems, and methods ofmaking the same, are presented herein according to embodiments of theinvention.

According to various embodiments, these metallic systems may include:pure copper (Cu), pure iron (Fe), or pure titanium (Ti), eachrepresentative of a specific crystallographic class, face centered cubic(FCC), body centered cubic (BCC), and hexagonal close packed (HCP),respectively. The untreated coarse-grained metal, upon being subjectedto cryogenic SMAT, will possess a gradient microstructure from thesurface to the bulk. That is, the as-processed metal will consist of anano- to submicrometer length scale grain structure at its surface,monotonically reaching, some distance away below the surface, theinitial, untreated grain size of the bulk. The ability to control thedepth below the surface and the effectiveness of this grain sizereduction is one the novel aspects of this invention.

According to various embodiments, these metallic systems may alsoinclude: multiple elements, wherein upon SMAT, the resultant componentelements may be in a discrete or disordered or alternatively ordered,layered, in a stratified arrangement, consisting of a single to multiplelayers, each with its own distinct features, displaying various levelsof intermixing between the initial constituents. Alternatively,adjustment of the conditions may generate a cellular structure on amultiscale level, with varying dimensions of the cell size.

In the various embodiments of the metallic systems, the initial metalmay have an untreated grain size greater than 20 to 50 μm. The variousexemplary metal systems, subjected to room temperature SMAT, may havegrain sizes as fine as 150 nm up to 360 nm. However, the grain sizereduction, only after one hour treatment, across the three exemplarymetallic systems, is considerably more effective at cryogenictemperatures: 60 nm up to 140 nm, or 45% to 62%, respectively.

For the purposes of this invention, various ASTM measurement standardsor their accepted equivalents were used. For example, for thedetermination of microhardness, ASTM E384-11e1, Standard Test Method forKnoop and Vickers Hardness of Materials, ASTM International, WestConshohocken, Pa., 2011 was used. Likewise, for the determination ofgrain size, the linear intercept method, quite common in microstructuralanalysis practices was used as well as the ASTM E112-13, Standard TestMethods for Determining Average Grain Size, ASTM International, WestConshohocken, Pa., 2013 was used as reference.

Furthermore, in these various embodiments, the depth of nanostructuring,as measured from the surface, can vary from 250 to 1000 μm. With the useof cryogenic temperatures, the depth changes, by as much as a factor oftwo, as determined by the crystallographic class of and the ease ofdeformation within the material. As such, for the BCC and HCP cases, thedepth of the affected region is significantly narrower, whereas, for theFCC case, it is wider. While it is preferred that a low temperature isused, more preferably the lowest attainable with the cooling system,various embodiments may necessitate altering and increasing theprocessing temperature to prevent the processed material fromprematurely fracturing or failing, due to it being too brittle or belowits ductile-brittle transition temperature.

These embodiments thus provide a new class of nanostructured andnanocrystalline metals and alloys which have significantly improvedproperties compared to their coarse grained counterparts. Moreover, theexemplary copper (Cu) metallic system subjected to room temperature SMATmay have a Vickers microhardness (HV) of about 110 kg/mm² at itssurface, while an equivalent sample, processed at a cryogenictemperature, may have an HV of about 155 kg/mm²; a significant increase.Similarly, the exemplary iron (Fe) metallic system subjected to roomtemperature SMAT may have an HV of about 220 to 250 kg/mm²; for iron(Fe) the resultant hardness difference between room and cryogenictemperature processed samples is not notable. Lastly, the exemplarytitanium (Ti) metallic system, processed with room temperature SMAT mayhave an HV of about 280 kg/mm²; whereas, the cryogenically treatedsample may have an HV of about 320 kg/mm². Again, in the spirit of thisinvention, these aforementioned microhardness values are illustrationsof typical properties for a specific processing condition, one hour ofSMAT. It is conceivable that longer SMAT times will lead to furtherhardness increases, especially at the surface. It is also believed thatshorter SMAT times, i.e. impact times, may be sufficient to achieve thedesirable improvement in final properties. Thus, in certain embodiments,the processing time, i.e. the total amount of time that the metal partsor parts are impacted with metal fragments (t_(i)) at reducedtemperature (T_(r)) is at least about 5 minutes, more preferably atleast about 10 minutes, still more preferably at least about 20 minutes,still even more preferably at least about 40 minutes and in someinstances at least about 1 hour or 60 minutes.

According to further embodiments, the SMAT process for forming theexemplary nanostructured Cu, Fe, or Ti metallic systems, comprised of agradient microstructure from nano- to micro- to macroscales, impartsthese metals with a preferred grain orientation or texture. For eachsystem, the texture evolves, as determined by the ease of deformationand grain reorientation when subjected to high strain and high stressconditions.

During SMAT, using a high-energy milling device may utilize a mixingvial to contain the metallic plate in addition to or replacing thevial's lid and, in some circumstances, may utilize a plurality ofmilling balls for inclusion within the mixing vial. In otherembodiments, especially for those when surface alloying is desired, inaddition to the plurality of milling balls, the vial may contain thepowder(s) to be alloyed therein. In certain circumstances, the vial andsteel milling media may be precoated with a specific metal topreferentially augment or reduce the presence of a specific element orspecies. The ball-to-powder mass ratio utilized by the high-energymilling device may be 1:1, more specifically, 10:1, or more.Furthermore, the milling balls may be comprised of high strength steelor ceramic. During the high-energy milling process, the vial and itscontents may be cooled to sub-ambient temperatures, more preferably, acryogenic temperature. This may be accomplished by cooling the millingdevice with any type of cryogenic liquid, more specifically, liquidnitrogen. Alternatively, the high-energy milling process may beperformed at ambient or room temperature. The high-energy millingprocess may be further improved using an additive or a surfactant. Insome instances, the vial and its contents may be continuously orsemi-continuously cooled during the high-energy milling process. Infurther embodiments, at the conclusion of the milling process, themetallic powder may be subjected to annealing by exposing it to elevatedtemperature in the range of about 300 to 800° C.

These embodiments thus provide a methodology for forming a new class ofnanostructured and nanocrystalline metals or composites which possessenhanced mechanical properties with a gradient grain sizemicrostructure.

These and other, further embodiments of the invention are described inmore detail, below.

BRIEF DESCRIPTION OF THE DRAWINGS

So that the manner in which the above recited features of the presentinvention can be understood in detail, a more particular description ofthe invention, briefly summarized above, may be had by reference toembodiments, some of which are illustrated in the appended drawings. Itis to be noted, however, that the appended drawings illustrate onlytypical embodiments of this invention and are therefore not to beconsidered limiting of its scope, for the invention may admit to otherequally effective embodiments, including less effective but also lessexpensive embodiments which for some applications may be preferred whenfunds are limited. These embodiments are intended to be included withinthe following description and protected by the accompanying claims.

FIG. 1 shows scanning electron micrographs of the gradient nature of thesubsurface grain size reduction in pure copper (Cu), comparing samplesprocessed at ambient and cryogenic temperatures, respectively.

FIG. 2 shows the x-ray diffraction patterns of the as-received, ambienttemperature processed, and cryogenically processed copper (Cu) samples,respectively.

FIG. 3 shows scanning electron micrographs of the gradient nature of thesubsurface grain size reduction in pure iron (Fe), comparing samplesprocessed at ambient and cryogenic temperatures, respectively.

FIG. 4 shows the x-ray diffraction patterns of the as-received, ambienttemperature processed, and cryogenically processed iron (Fe) samples,respectively.

FIG. 5 shows scanning electron micrographs of the gradient nature of thesubsurface grain size reduction in pure titanium (Ti), comparing samplesprocessed at ambient and cryogenic temperatures, respectively.

FIG. 6 shows a high magnification scanning electron micrographs forcopper (Cu), iron (Fe), and titanium (Ti) metal types revealing thedramatic decrease in grain size with the use of cryogenic processingconditions versus those at ambient conditions.

FIG. 7 shows the x-ray diffraction patterns of the as-received, ambienttemperature processed, and cryogenically processed titanium (Ti)samples, respectively.

FIG. 8 shows the Vickers microhardness values of the copper (Cu), iron(Fe), and titanium (Ti) samples, as a function of depth away from thesurface for both the ambient and cryogenic temperature processedsamples.

DETAILED DESCRIPTION OF INVENTION

Unlike that in the prior art, we utilize for the first time a novel lowtemperature or cryogenic SMAT process on commercially available,oxygen-free high-conductivity (OFHC) copper (Cu), pure iron (Fe), andpure titanium (Ti) and show the advantages in terms of microstructurerefinement as compared to conventional ambient or room temperature (RT)SMAT processing. While SMAT processes in a conventional configurationhave been applied to a number of metallic systems, including Cu, Ni, Ti,Fe, and stainless steels, there currently is very limited datademonstrating decoupling of thermomechanical effects during theprocessing, especially at cryogenic temperatures (e.g., below −150° C.)where dynamic recovery and recrystallization are dramaticallysuppressed. By utilizing equipment developed for high-energy cryogenicmechanical alloying or milling, which is not typically used, it ispossible to show the substantial microstructure refinement of thecryogenic SMAT process in OFHC Cu, pure Fe, or pure Ti over what isattainable at RT.

Thus, in certain preferred embodiments the present invention provides amethod of modifying the surface of a metal part that includes: providingat least one metal part that is formed from a first metal composition;providing a plurality of metal fragments that are formed from a secondmetal composition wherein said metal fragments have a size that issignificantly smaller than the size of the at least one metal part;reducing the temperature of the at least one metal part and theplurality of metal fragments; impacting the at least one metal part withthe plurality of fragments at a reduced temperature (T_(r)), for animpact processing time (t_(i)); wherein, the at least one metal part issubjected to bombardment by the plurality of metal fragments at areduced temperature (T_(r)), resulting in the as-received grain size ofthe at least one metal part to be reduced, by several orders ofmagnitude, and possessing a gradient structure from the impact surfaceof the at least one metal part into the interior of the bulk of the atleast one metal part.

In the exemplary embodiments provided below, the metal fragments weresteel balls or milling balls. However, ceramic milling balls may be usedor some other material exhibiting properties similar in nature.Suggested steel milling balls will have a diameter of at least about0.015 mm and more preferably at least about 1.5 mm and may have adiameter as large as about 2.5 cm. In certain embodiments, the millingballs are formed from tungsten. Thus, the mass of an individual millingball may range from about 1.01 gram to about 160 grams. Desirably, theproperties of the milling balls, including but not limited to hardness,should be similar to or greater than the properties of the materialbeing milled or SMATed. However, in certain embodiments the millingballs may be abraded during the process to produce an alloy or compositestructure at the surface of the part or material that is being milled orSMATed. In certain other embodiments a powder of is included in themilling container to produce an alloyed surface or to introduce a secondphase at the SMATed surface.

In the exemplary embodiments provided below, the at least one metal partincluded a plurality of metal plates. However, in commercialapplications it is suggested that the at least one metal part is afinished part or an almost finished part, for example a head of a golfclub, a gear or gear tooth or a part of a gun or an engine componentprior to assembly or packaging. The composition of the metal part(s) maybe same or different from the composition of the metal fragments. In theexemplary embodiments provided below, the composition of the metalparts, e.g. metal plates, and the composition of the fragments weredifferent. However, in some applications it is may be desirable that thecomposition of the metal part(s) may be same as the composition of themetal fragments or substantially the same as the composition of themetal fragments. For example, both the part(s) and the impact fragmentsmay be formed from high-conductivity (OFHC) copper. In certainembodiments the surface at least one part that is being SMATed atreduced temperature is cut away and then used as a sub part orcomponent. Additionally, the SMATed surface may be further processed.For example, the high density of grain boundaries in a nanocrystalllinemicrostructure also increases the available diffusion pathways at thesurface for secondary processing such as carburization or nitriding,enabling decreases in process temperature of possibly improved speciesabsorption.

Samples were prepared by cutting commercially available OFHC Cu, pureFe, or pure Ti into disks 6.35 cm in diameter and 0.6 cm thick. Thesesamples were then polished to a mirror finish. Nominally supplied highenergy SPEX® SamplePrep Corporation mill vial lids were then replaced bythe samples and sealed in a high-purity argon glovebox. In each case,the hardened steel vials were loaded with 17 5/16″ diameter and 16¼″diameter 440 C stainless steel ball bearings to constitute a total massof 50 g. For the Cu and Ti samples, the vial was precoated with a thinlayer of Cu or Ti, respectively. This was to limit potentialcontamination of the as SMATed surface from the Fe constituent of themilling vial. The precoating was accomplished by placing approximately0.5 g of each respective powder into the vial and operating the SPEXmill for 10 min at room temperature. The actual SMAT samples were eitherprocessed at cryogenic or RT for a period of 1 hour. For the purposes ofthe invention, low or cryogenic temperatures are defined as temperatureswell below ambient room temperature conditions, preferably below about−50° C. (223K), not greater than about −100° C. (173K), more preferablynot greater than about −150° C. (123K), and most preferably not greaterthan about −196° C. (77K). RT milling was accomplished by loading thevials into a commercially available single vial SPEX SamplePrep Model8000M mill, while cryogenic milling was performed in a modified SPEXSamplePrep Model 8000M mill. The modified mill was equipped with aTeflon sleeve into which the sealed steel vials could be inserted. TheTeflon sleeve was fitted to allow the inflow and outflow of liquidnitrogen at a temperature of −196° C. (77K) to envelope the outside ofthe steel vial. After the cryogenic and RT SMAT processing, the sampleswere mounted and polished for microstructure analysis.

The microstructure was analyzed using optical microscopy, X-raydiffraction for grain size and strain analysis, and scanning electronmicroscopy (SEM) for chemical, grain size and orientation analysis.Electron backscatter imaging in the SEM was used to analyze the grainsize. Optical imaging of etched samples was used to assess the specificgrain size and defect morphology. A dual-beam FEI Nano600 FIB was usedto prepare samples for electron backscatter diffraction (EBSD) imagingand produce ion-channeling contrast images to highlight the activedeformation mechanisms during cryogenic SMAT processing.

FIGS. 1a and 1b are SEM images of the etched sample cross-sections oftwo copper (Cu) samples prepared by RT and cryogenic SMAT processing for1 hour, respectively. For the purposes of the invention, the duration ofprocessing is defined as times being at least 5 minutes, preferably atleast 10 minutes, more preferably at least 30 minutes, and mostpreferably at least 60 minutes. Note, however, the duration ofprocessing is a strong function of the ductile-brittle temperature ormalleability of the metal being treated, its crystallographic class, andthe temperature dependence of its malleability. As such, for thepurposes of the invention, there is no upper bound in processing times,except that determined by the failure of the metal plate due to voidformation, fracture, spallation, and subsequent fragmentation. As such,under certain conditions, processing times as long as two to three hoursmay be feasible. The exemplary and representative micrographs illustratethe change in grain structure as a function of depth from the SMATsurface, which is analogous to how the microstructure evolves as afunction of time for the SMAT process. The SMAT surface had up to 20 μmnon-continuous layer with some Fe contamination from the use ofstainless steel ball bearings; this is present in both RT and cryogenicsamples.

There is a noticeable transition from nanostructured orultrafine-grained grain structure to a region of banded structures tothe bulk microstructure within the first 500 μm. The back surface of theSMAT specimens (images to the left, denoted by a square) is 0.6 cm fromthe SMAT surface. The microstructure near the back surface represents apristine microstructure unaffected by the SMAT process; the averageequiaxed grain size in this region is 150 μm in diameter with sometwinning present. The average grain size for the cryogenic sample (140nm mean diameter) is approximately 60% smaller than the grain sizeachieved through RT processing (355 nm mean diameter).

The cryogenic SMAT sample maintains a constant finer grain size to amuch deeper penetration depth. That is, in contrast to the cryogenicsample, the RT processed sample appears to have a continuous monotonicincrease in grain size as a function of depth into the plate.Furthermore, the overall region of grain refinement, defined here as thevolume containing grains smaller than 10 μm in diameter, is found tovary with processing conditions. This region is approximately 150 μmdeep for the cryogenically processed sample and 300 μm deep for the RTprocessed sample, respectively.

As shown in FIG. 2, X-ray diffraction patterns were collected from theas-received OFHC copper (Cu) plate, RT, and cryogenic SMAT samples. Uponinspection for the fundamental Bragg crystallographic reflections, thefull widths at half-maximum increased and the amplitude decreased in thefollowing order: as-received, 1 hour RT condition and 1 hour cryogeniccondition. These trends are consistent with a decrease in grain sizeand/or an increase in local strain with processing conditions. TheScherrer estimates of grain size indicate that the average grain sizefor the cryogenic SMAT samples is lower than that calculated for the RTSMAT samples. However, the grain size estimates using the Scherrerformula were significantly lower than the previously measured grain sizeusing SEM micrographs.

The texture change during SMAT can also be measured by the change in therelative intensity of the fundamental crystallographic reflections ofCu. For the cryogenic SMAT sample, the relative intensities of thefundamental reflections of pure Cu, as given by the JCPDS index, are:(111) 100%, (200) 46%, (220) 20%, (311) 17% and (222) 5%. For theunprocessed plate, the second and third reflections have higher percentrelative intensities: (200) 58% and (220) 63%. This is in contrast tothe texture generated during the SMAT process. Hence, in general, SMATinduces a texture wherein the (111) orientation is favored relative tothe other reflections. Furthermore, cryogenic SMAT processing enhancesthis type of texturing over RT SMAT processing. For the cryogenicallyprocessed sample, the ratio of intensities of the (200) and (111)reflections increases to a slightly higher value than that with RTprocessing.

FIGS. 3a and b show the equivalent SEM images of the etched iron (Fe)sample cross-sections processed for 1 hour, respectively. In theseimages the contrast in grain morphology between the RT and cryogenicprocessed samples is rather striking. For iron (Fe) processed at RT, thedeformed and reduced grains show a significant texturing effect, whereinthe grains appear in layers and are stratified perpendicular to theplate normal. Whereas, the iron (Fe) processed at a cryogenictemperature, the as-deformed region is less stratified and moreequiaxed. Comparing the affected depths between conditions, the RTprocessed specimen is significantly deeper than that processed atcryogenic temperature. For the former, the highly refined grain sizeregion is about 200 μm deep; for the latter, the corresponding region isonly about 50 μm. An intermediate region between the unaltered bulk isalso wider for the sample processed at RT.

Unlike the FCC copper (Cu) samples which showed a dramatic increase inthe (111) crystallographic orientation relative to all of the others,the BCC iron (Fe) samples show a considerably lesser effect. As shown inFIG. 4, are the X-ray diffraction patterns of the as-received iron (Fe)plate, RT, and cryogenic SMAT samples. Inspection of the fundamentalBragg reflections show that while the full widths at half-maximum didsignificantly change, the amplitudes did not decrease between theconditions presented. This effect is directly related to peak broadeningdue to grain size reduction, however, without, a significant change incrystallographic orientations. That is, aside from a reduction in grainsize and pancaking of the grains, SMAT, in either RT or cryogenicconditions, does not introduce grain reorientation.

As was seen for both copper (Cu) and iron (Fe), there is a correspondingdecrease in the size of region of nano to ultrafine-scale grains for HCPtitanium (Ti) as well. The grains appear to be mostly equiaxed; absentis the layering or stratification of the grains. In FIGS. 5a and 5 b,SEM images of the etched sample cross-sections of two titanium (Ti)samples prepared by RT and cryogenic SMAT processing for 1 hour,respectively, show that submicrometer grains persist to a depth of about75 μm in the RT processed sample, whereas, this depth is only about 40μm for the cryogenic sample. Again, the transition region from thenano-scale to macroscale grains is deeper for the cryogenic condition.However, overall, the depths of the affected regions are about the same;450 μm. This effect is similar to the case of copper (Cu), but differentfrom that of iron (Fe).

X-ray diffraction patterns of the as-received titanium (Ti) plate, RT,and cryogenic SMAT samples are shown in FIG. 6. The full widths athalf-maximum of the Bragg peaks illustrate significant line broadeningfor both SMAT conditions. In fact, the two primary reflections of (002)and (101), partially and completely overlap for the RT and cryogenicSMAT conditions, respectively. Note, with the exception of the growth ofthe (101) peak, the there is little change in the relative peak heightsof the other peaks between the processing conditions. Given the factthat HCP titanium (Ti) has a limited number of operative deformationslip systems it is most likely, that recrystallization is the primarygrain refinement mechanism in this material. This is consistent with theobserved grain morphology and lack of texturing as indicated by thechanges in relative peak heights.

For the purposes of the invention and the preferred embodimentsdescribed herein, it is important to realize that, whereas, the one hourprocessing time may have been closer to optimum conditions for FCCcopper (Cu), however, this processing time may not have been the casefor latter, namely the BCC iron (Fe) and HCP titanium (Ti) embodiments.That is, it is implied from a comparison of results that the resultanttexturing in the latter systems have not yet fully evolved. In otherwords, different crystallographic systems will develop differently dueto their intrinsic properties and underlying deformation mechanisms.Thus, it is likely that longer processing times would have resulted inthe evolution of a stronger texture in the other exemplary metals.

Another factor in these embodiments is the evolution of steady stateequilibrium conditions during SMAT processing, wherein, the heatgeneration due to deformation is offset by active cooling. That is, inall likelihood, the processing temperature, while fixed at −196° C.(77K), primarily, for convenience, delivery, and availability of liquidnitrogen will have a significant effect on the effectiveness of thegrain size refinement process. Thus, for the purposes of this invention,it is hypothesized that it is highly likely that at lower or highertemperatures, corresponding to alternate equilibrium conditions, wouldresult in a different more favorable outcome, i.e., potentially finergrain size reduction for the latter metal systems.

FIG. 7 reveals higher magnification SEM images of regions located atgreater depths below the SMAT surface. There are major microstructuredifferences with respect to the processing conditions. First, thecryogenically processed samples maintain a near-equiaxed grainmorphology, whereas the RT processed samples have larger regions wherethe grain morphology is distorted from its initial equiaxed shape.Second, a high density of etch pits, mostly likely associated withdislocations intersecting the polished surface, are present in the RTsample. When metals undergo severe plastic deformation at or near RT,dislocation slip and deformation twinning are the principal modes ofdeformation. While both FCC and BCC metals have adequate numbers ofoperational slip systems, there is a considerable limitation of suchslip systems in HCP metals. Regardless, dislocation tangles and theirspecific subgrain structures (dislocation cells, walls, geometricallynecessary boundaries and incidental dislocation boundaries) aregenerated and equilibrated by thermally activated processes. As such,these are the means for achieving grain refinement during SMATprocessing.

The cryogenic SMAT samples, however, show very different microstructureevolution to that of the RT SMAT samples. At lower temperatures, thedislocation-based processes are suppressed, as indicated by the lack ofetch pits in the cryogenically processed samples. Moreover, in contrastto the RT SMAT samples, a large number of banded structures wereobserved within the large grains of the cryogenic SMAT samples. In manycases, these banded structures initiate from the grain boundaries andeither terminate at the inside wall of the same grain or intersect theopposite side boundary. The spacing between many of the bands is lessthan 10 μm. In general, these bands contain small ultrafine equiaxedgrain structures or structures that are elongated and parallel to eachother across the width of the band. These observations are consistentwith observations of twin/matrix bundles, bamboo nanograins, and shearbands that evolve as a function of strain during dynamic plasticdeformation under liquid nitrogen temperatures. These observationsindicate a shift in the dominant deformation mechanism during SMAT fromdislocation-mediated behavior at room temperature to twinning/shearband-mediated at cryogenic temperatures.

A time evolution of the microstructure can be constructed for thecryogenic SMAT process. First, due to the high-strain-rate impacts atcryogenic temperatures, twin/matrix bundles and/or shear bands originateat the surface or at internal stress concentration sites (grainboundaries, triple junctions, etc.). The repeated impacts generateoverlapping twin/matrix bundles and/or shear bands, which effectivelyrefine the large micrometer-sized grains. Additional grain refinementmay occur within any shear band present due to dynamicrecrystallization. As the internal stress continues to accumulate, themicrostructure evolves by producing ultrafine grains with long aspectratios. The accumulated shear strain leads to fragmentation and rotationof these elongated grains, resulting in an equiaxed nanocrystallinemicrostructure.

FIG. 8 displays the results of microhardness measurements taken from thethree metal specimen types both after being subjected to RT or cryogenicSMAT. The three embodiment exhibit a monotonic fall off in microhardnessthat strongly depends on the local grain size at the relative depthmeasured from the surface. Whereas, the RT induced hardness change isgradual from the surface to the interior, for samples processed atcryogenic temperatures, this hardness change is more rapid, presumablydue to the shallower nature of the structurally modified region. Themost significant change occurs for pure copper (Cu), then pure titanium(Ti), and lastly pure iron (Fe). The figure clearly illustrates thevariations between the three embodiments. The advantage for the use ofcryogenic conditions is most beneficial in the case of the softer copper(Cu), however, if a sharper hardness profile and grain sizerestructuring is desired, this could be achieved by altering the SMATconditions, with lengthening the duration and potentially lowering thetemperature further. The embodiment using titanium (Ti) illustrates thatquite well. Note, in the case of pure iron (Fe), there is little to nobenefit to the use of cryogenic processing conditions, below thesurface. However, this does not preclude the use and the resultantbenefits that derive from lower temperatures or longer SMAT times.

Supporting and/or additional details may be found in a journal articletitled “Enhancing grain refinement in polycrystalline materials usingsurface mechanical attrition treatment at cryogenic temperatures”Scripta Materialia 69 (2013) 461-464 by Dr. Kristopher A. Darling et al.which is hereby incorporated by reference herein.

The foregoing description, for the purposes of providing an explanation,has been described with reference to specific embodiments. However, theillustrative discussions above are not intended to be exhaustive or tolimit the invention to the precise forms and conditions disclosed. Manymodifications and variations are possible in view of the aboveteachings. The embodiments were chosen and described in order to bestexplain the operating principles of the present disclosure and itspractical applications to thereby enable others skilled in the art tobest utilize the invention and its various embodiments with variousmodifications as may be suited to the particular use contemplated.

While the foregoing is directed to embodiments of the present invention,other and further embodiments of the invention may be devised withoutdeparting from the basic scope thereof, and the scope thereof isdetermined by the claims that follow.

1. A surface mechanical attrition treatment process for a metal partcomprising: providing at least one metal part that is formed from afirst metal composition; providing a plurality of metal fragments thatare formed from a second metal composition wherein said metal fragmentshave a size that is significantly smaller than the size of the at leastone metal part; reducing the temperature of the at least one metal partand the plurality of metal fragments; impacting the at least one metalpart with the plurality of fragments at a reduced temperature (T_(r)),for an impact processing time (t_(i)); wherein, the at least one metalpart is subjected to bombardment by the plurality of metal fragments ata reduced temperature (T_(r)), resulting in the as-received grain sizeof the at least one metal part to be reduced, by several orders ofmagnitude, and possessing a gradient structure from the impact surfaceof the at least one metal part into the interior of the bulk of the atleast one metal part.
 2. The surface mechanical attrition treatmentprocess of claim 1, wherein the reduced temperature (T_(r)) is notgreater than about −50° C.
 3. The surface mechanical attrition treatmentprocess of claim 1, wherein the reduced temperature (T_(r)) is notgreater than about −100° C.
 4. The surface mechanical attritiontreatment process of claim 1, wherein the reduced temperature (T_(r)) isnot greater than about −150° C.
 5. The surface mechanical attritiontreatment process of claim 1, wherein the reduced temperature (T_(r)) isnot greater than about −196° C.
 6. The surface mechanical attritiontreatment process of claim 1, wherein the refined grain size of themetallic system at its surface is about 1000 nm or less.
 7. The surfacemechanical attrition treatment process of claim 3, wherein the roomtemperature refined grain size of the metallic system at its surface isabout 355 nm or less.
 8. The surface mechanical attrition treatmentprocess of claim 3, wherein the cryogenic refined grain size of themetallic system at its surface is about 140 nm or less.
 9. The surfacemechanical attrition treatment process of claim 3, wherein the roomtemperature processed metallic system has a surface Vickersmicrohardness of about 110 kg/mm² or more at room temperature.
 10. Thesurface mechanical attrition treatment process of claim 3, wherein thecryogenic temperature processed metallic system has a surface Vickersmicrohardness of about 150 kg/mm² or more at cryogenic temperature. 11.The surface mechanical attrition treatment process of claim 4, whereinthe room temperature refined grain size of the metallic system at itssurface is about 635 nm or less.
 12. The surface mechanical attritiontreatment process of claim 4, wherein the cryogenic refined grain sizeof the metallic system at its surface is about 350 nm or less.
 13. Thesurface mechanical attrition treatment process of claim 4, wherein theroom temperature processed metallic system has a surface Vickersmicrohardness of about 230 kg/mm² or more at room temperature.
 14. Thesurface mechanical attrition treatment process of claim 4, wherein thecryogenic temperature processed metallic system has a surface Vickersmicrohardness of about 260 kg/mm² or more at room temperature.
 15. Thesurface mechanical attrition treatment process of claim 5, wherein theroom temperature refined grain size of the metallic system at itssurface is about 155 nm or less.
 16. The surface mechanical attritiontreatment process of claim 5, wherein the cryogenic refined grain sizeof the metallic system at its surface is about 60 nm or less.
 17. Thesurface mechanical attrition treatment process of claim 5, wherein theroom temperature processed metallic system has a surface Vickersmicrohardness of about 280 kg/mm² or more at room temperature.
 18. Thesurface mechanical attrition treatment process of claim 5, wherein thecryogenic temperature processed metallic system has a surface Vickersmicrohardness of about 320 kg/mm² or more at room temperature.
 19. Thesurface mechanical attrition treatment process of claim 1, wherein theimpact processing time (t_(i)) is at least about 5 minutes.
 20. Thesurface mechanical attrition treatment process of claim 1, wherein theimpact processing time (t_(i)) is at least about 10 minutes.
 21. Thesurface mechanical attrition treatment process of claim 1, wherein theimpact processing time (t_(i)) is at least about 20 minutes.
 22. Thesurface mechanical attrition treatment process of claim 1, wherein theimpact processing time (t_(i)) is at least about 40 minutes.
 23. Thesurface mechanical attrition treatment process of claim 1, wherein theimpact processing time (t_(i)) is at least about 60 minutes.
 24. Amethod of modifying the surface of a metal part, the method comprisingthe steps of: providing at least one metal part t; providing a pluralityof metal fragments wherein said metal fragments have a size that issignificantly smaller than the size of the at least one metal part;reducing the temperature of the at least one metal part and theplurality of metal fragments; and impacting the at least one metal partwith the plurality of fragments at a reduced temperature (T_(r)), for animpact processing time (t_(i)).